Optimization of the Growth Process of Double Perovskite Pr2−δNi1−xMn1+xO6−y Epitaxial Thin Films by RF Sputtering

Epitaxial thin films of Pr2−δNi1−xMn1+xO6−y (PNMO) double perovskite were grown on (001)-oriented SrTiO3 substrates by RF magnetron sputtering. The influence of the growth parameters (oxygen pressure, substrate temperature, and annealing treatments) on the structural, magnetic and transport properties, and stoichiometry of the films was thoroughly investigated. It is found that high-quality epitaxial, insulating, and ferromagnetic PNMO thin films can only be obtained in a narrow deposition parameter window. It is shown that a careful selection of the growth conditions allows for obtaining a high degree of Ni/Mn cation ordering, which is reflected in the values of the Curie temperature, TC, and saturation magnetization, MS, which are very close to those of bulk material.


Introduction
Spintronics has been pointed out as a potential alternative technology for the development of multifunctional electronics combining data storage and transmission, as well as logical operations, with improved energy efficiency. The generation and control of pure spin currents [1] or highly polarized electron currents [2] is at the cornerstone of spintronics. One of the most common ways for generating a highly polarized electron current is the spin filtering through the ferromagnetic-insulator-ferromagnetic tunnel barriers [3]. In addition, the same can be achieved by using a single element, a ferromagnetic-insulating (FM-I) barrier [4,5]. However, FM-I materials are very scarce. In this regard, double perovskite (DP) oxides of the R 2 NiMnO 6 family (RNMO, where R is a rare earth element) have received considerable attention for presenting this behavior. Their multifunctional physical properties make them potential candidates for technological applications beyond spintronic devices, such as magnetodielectric capacitors and multiple state logic devices [6][7][8][9][10]. In this family, La 2 NiMnO 6 (LNMO; R = La) has attracted much more attention for its near-room-temperature ferromagnetic transition (T C ≈ 280 K), insulating characteristics, magneto-dielectric effect, and spin-phonon coupling [7,11,12]. The ferromagnetic (FM) ordering in RNMO systems is explained in terms of the superexchange interaction between Ni 2+ (d 8 , t 6 2g e 2 g , S = 1) and Mn 4+ d 3 , t 3 2g e 0 g , S = 3/2 , which is FM in 180 • geometry according to the Goodenough-Kanamori rules [7,[13][14][15]. For this reason, the magnetic properties of these materials are found to be very sensitive to the Ni/Mn cationic ordering inside of the DP structure.
Previous studies on RNMO systems have demonstrated that the substitution of La 3+ by another rare earth element, R 3+ , with a smaller ionic radius induces significant modifications in the Ni-O-Mn bond angle directly involved in the superexchange mecha-

Experimental Details
The PNMO thin films were deposited on (001)-oriented STO substrates by the RF magnetron sputtering technique by using a stoichiometric Pr 2 NiMnO 6 target prepared by the solid-state reaction method. Before the deposition, the substrates were cleaned in an ultrasonic bath with Milli-Q water and then annealed at 1000 • C in air for 2 h to obtain a clean and smooth TiO 2 -terminated surface [29]. Stoichiometric target (ϕ = 1.3") was synthesized by mixing the primary oxide powders (99.99% pure), i.e., La 2 O 3 , Pr 2 O 3 , MnO 2 , and NiO. At the end, the pressed pellet was thermally treated up to 1400 • C for 48 h in air, as this process ensures to achieve a homogeneous and dense target. The X-ray powder diffraction pattern analysis of the final target of Pr 2 NiMnO 6 with no additional spurious phase present is shown in Figure S1 in Supporting Information. The XRD peaks were indexed on the basis of monoclinic P 2 1 /n (single cell) according to the Rietveld refinement. The average of the lattice parameters of the pseudocubic phase of Pr 2 NiMnO 6 was 3.871 Å.
Films were grown at different background oxygen pressures (100-750 mTorr) and growth temperatures ranging from 200 to 900 • C. After thin film growth, the films were annealed in situ at the same growth temperature for 1h under 420 Torr O 2 and then slowly cooled down to room temperature at 10 • C/min. All PNMO films were prepared at a fixed RF power (40 W) and a fixed target-to-substrate distance (5 cm).
The chemical composition and stoichiometry ratio of the PNMO thin films were determined by wavelength dispersive spectrometer (WDS) electron probe microanalysis (EPMA) using a (CAMECA, Gennevilliers, France) SX-50 electron microprobe equipped with four wavelength-dispersive X-ray spectrometers. The electrical transport properties were measured in a Physical Properties Measurement System (PPMS, Quantum Design, San Diego, CA, USA) by the four-probe technique. Magnetization measurements were done CA a superconducting quantum interferometer device (SQUID, Quantum Design).

Oxygen Pressure Deposition
First, the optimal deposition conditions for growing the ordered double perovskite PNMO thin films on SrTiO 3 substrates were examined. We started the optimization of the growth parameters using high growth temperatures (850 • C), following previous studies [26]. Under these conditions, a series of PNMO thin films were grown for 1 h under different oxygen pressures (ranging from 100 to 750 mTorr) by keeping fixed the growth temperature (T = 850 • C). Figure S2a in Supporting Information shows the XRR measurements that were carried out to determine the thickness of the samples. Thicknesses were obtained from the positions of the minima in the XRR curve. In Figure S2b in Supporting Information, one can appreciate that when the oxygen pressure increases from 100 to 750 mTorr, the thickness decreases monotonously from 95 nm to 28.5 nm. Accordingly, the growth rate decreases from 1.6 to 0.5 nm/min, respectively. Normally, by increasing the oxygen pressure in the chamber, the plasma becomes broader, and the mean free path of the atoms decreases, thus decreasing the quantity of the arriving adatoms and the growth rate.
The crystallinity quality of the films has been evaluated by X-ray diffraction. Figure 1a shows θ/2θ scans of the PNMO thin films at a fixed deposition temperature of 850 • C under various oxygen pressures. On all the XRD patterns, the (00l) reflections of the PNMO phase (indexed in the pseudocubic notation), which take place at 2θ values slightly higher than the ones of SrTiO 3 , can be clearly observed (see dashed vertical red lines), thus indicating a good crystalline quality of the films. For closer inspection, Figure 1b shows the (002) reflection of both the STO substrate and the PNMO film for each sample. One can observe that the position of the (002) peak from the PNMO on the STO occurs at higher 2θ angles than the bulk PNMO target value of 3.871 Å located at 2θ = 46.9 • (see dashed vertical red line in Figure 1b). This indicates a shrinking of the out-of-plane c lattice parameter, in agreement with the lattice mismatch imposing a tensile in-plane strain. In Figure 1c, the estimated out-of-plane c lattice parameter is plotted. At 100 mTorr O 2 , the PNMO (002) reflection is located at 2θ ≈ 47.17 • , corresponding to an out-of-plane cell parameter of 3.864 Å. By increasing the oxygen pressure from 100 to 300 mTorr, the PNMO (002) peak shows a shift towards the slightly higher 2θ angles (see arrow) if compared with its bulk position. This displacement is associated with a decrease of the out-of-plane c lattice parameter (see Figure 1c). At intermediate oxygen pressures (350 and 400 mTorr), a shift of the PNMO peak towards lower angles can be seen (see arrow), indicating an expansion of the out-of-plane c lattice parameter of the PNMO film, approaching the bulk PNMO value of 3.871 Å (see Figure 1c). With the further increasing of the oxygen pressure (above 400 mTorr), the out-of-plane c lattice parameter decreases again. The latter could be associated with a variation in the concentration of the defects and oxygen vacancies into the structure under different oxygen pressures.  Figure S1).
Finally, in addition to the PNMO reflections, the XRD patterns show additional reflections corresponding to the parasitic phases. In all the XRD patterns of the samples grown at very high and low oxygen pressures, the (002) reflection of the NiO phase can be clearly observed at 2θ ≈ 43.4° (see the asterisk and dashed blue line in Figure 1a,b). At the intermediate oxygen pressure (from 300 to 400 mTorr), a slight shift of this peak towards greater angles can be observed around 2θ ≈ 43.7°. This slight shift in the NiO peak could be attributed, at first glance, to a change in the oxygen content of the nickel oxide (NiOx). On the other hand, the reason for the presence of the NiO phase in the PNMO samples (grown/annealed at 850 °C) is not well understood, although this phase segregation may be of the same origin as seen in the LNMO thin films, as previously reported [30]. Moreover, it could be a consequence of a certain degree of Pr deficiency, as revealed by the electron probe microanalysis (EPMA) measurements (see below). This Pr deficiency may indeed induce the formation of secondary phases in the R2NiMnO6 double perovskite, as reported by Dass et al. [19]. Similarly, a secondary phase located at 43.24°, identified as NiO, was also observed in the XRD pattern of the Nd2NiMnO6 samples [31]. Therefore, according to the XRD data, the PNMO films grown at 850 °C under different oxygen pressures have a good crystalline quality. In addition, although the secondary phases are formed (NiO) together with the PNMO phase, the films grow epitaxial on the (001) STO substrates.
The surface morphology of the PNMO films was examined by AFM and SEM measurements (see Figure 2a,b and Figure S3 in Supporting Information). All the AFM images (left-hand images) show a granular morphology (no step-terrace morphology detected). Furthermore, in Figure 2c, one can appreciate that the PNMO films deposited under low oxygen pressures (ranging from 100 to 300 mTorr) have high root mean square (rms) roughness values (rms > 1.5 nm). By increasing the oxygen pressure from 350 to 750 mTorr, the film morphologies become flatter and smoother with low roughness values (rms ≈ 0.5 nm). At the same time, the presence of very fine grains distributed all over the film surface can be observed. The lower roughness could be related to the fact that the films grown under high oxygen pressure are thinner (see Figure S2b in Supporting Information). However, the decrease of the film thickness with oxygen pressure is monotonous, while the roughness decreases abruptly above 350 mTorr, indicating that the growth process at higher pressure favors a flatter surface.  Figure S1).
Finally, in addition to the PNMO reflections, the XRD patterns show additional reflections corresponding to the parasitic phases. In all the XRD patterns of the samples grown at very high and low oxygen pressures, the (002) reflection of the NiO phase can be clearly observed at 2θ ≈ 43.4 • (see the asterisk and dashed blue line in Figure 1a,b). At the intermediate oxygen pressure (from 300 to 400 mTorr), a slight shift of this peak towards greater angles can be observed around 2θ ≈ 43.7 • . This slight shift in the NiO peak could be attributed, at first glance, to a change in the oxygen content of the nickel oxide (NiO x ). On the other hand, the reason for the presence of the NiO phase in the PNMO samples (grown/annealed at 850 • C) is not well understood, although this phase segregation may be of the same origin as seen in the LNMO thin films, as previously reported [30]. Moreover, it could be a consequence of a certain degree of Pr deficiency, as revealed by the electron probe microanalysis (EPMA) measurements (see below). This Pr deficiency may indeed induce the formation of secondary phases in the R 2 NiMnO 6 double perovskite, as reported by Dass et al. [19]. Similarly, a secondary phase located at 43.24 • , identified as NiO, was also observed in the XRD pattern of the Nd 2 NiMnO 6 samples [31]. Therefore, according to the XRD data, the PNMO films grown at 850 • C under different oxygen pressures have a good crystalline quality. In addition, although the secondary phases are formed (NiO) together with the PNMO phase, the films grow epitaxial on the (001) STO substrates.
The surface morphology of the PNMO films was examined by AFM and SEM measurements (see Figure 2a,b and Figure S3 in Supporting Information). All the AFM images (left-hand images) show a granular morphology (no step-terrace morphology detected). Furthermore, in Figure 2c, one can appreciate that the PNMO films deposited under low oxygen pressures (ranging from 100 to 300 mTorr) have high root mean square (rms) roughness values (rms > 1.5 nm). By increasing the oxygen pressure from 350 to 750 mTorr, the film morphologies become flatter and smoother with low roughness values (rms ≈ 0.5 nm). At the same time, the presence of very fine grains distributed all over the film surface can be observed. The lower roughness could be related to the fact that the films grown under high oxygen pressure are thinner (see Figure S2b in Supporting Information). However, the decrease of the film thickness with oxygen pressure is monotonous, while the roughness decreases abruptly above 350 mTorr, indicating that the growth process at higher pressure favors a flatter surface. Additionally, the effect of the oxygen pressure on the film morphology was also explored by the SEM measurements. Figure 2a,b and Figure S3 in Supporting Information show the SEM images (right-hand images) of the corresponding AFM ones. Films deposited at low oxygen pressures (PO2 < 350 mTorr) show a rather flat morphology covered with small grains distributed homogeneously over the film surface. The latter corroborates the morphology observed by AFM and the increase of the roughness of the films. On the other hand, at high oxygen pressures (PO2 ≥ 350 mTorr), the SEM images reveal a flat and smooth morphology, although some small grains randomly distributed over the film surface can be observed. As mentioned before, high pressure favors a smoother surface. The presence of these small grains at the film surface may be due to the presence of the NiO secondary phase, as previously detected by the XRD patterns (see Figure 1a,b). On the other hand, the existence of a nanostructured surface was also detected in the LNMO systems, and it was attributed to the formation of NiOx aggregates after high annealing temperatures (800-900 °C) [30].

Deposition Temperature
According to the results on the PNMO films grown at different oxygen pressures, good ferromagnetic properties were obtained for the samples deposited at intermediate oxygen pressures ranging from 350 to 400 mTorr O2 (vide ultra). Therefore, in order to further optimize the deposition conditions, the PNMO thin films on the STO (001) substrates were grown at different substrate temperatures (ranging from 200 °C to 900 °C) by keeping a fixed oxygen pressure of 350 mTorr for 1h. Figure S4a in Supporting Information shows a typical XRR curve of a PNMO film deposited at high temperature (T = 700 °C). The thickness dependence as a function of the growth temperature is shown Figure S4b in Supporting Information. It can be observed that the film thickness remains almost constant above 600 °C. The thickness values found via the XRR analysis for all the PNMO samples grown at different temperatures are very similar (~48 nm), with a growth rate of about 0.8 nm/min. In a particular case, the sample deposited at a substrate temperature of 200 °C was thicker (~54 nm). Figure 3a shows the XRD θ/2θ scans of the ~48 nm-thick PNMO films grown under 350 mTorr O2 at different temperatures. In addition, in all the XRD patterns, only the (00l) reflections of the PNMO phase were clearly observed in the samples grown at temperatures above 600 °C (see the dashed vertical red lines), indicating that high growth temperatures favor a good crystalline quality. For more details, Figure 3b shows the (002) reflection of both the STO substrate and the PNMO film. As the growth tempera- Additionally, the effect of the oxygen pressure on the film morphology was also explored by the SEM measurements. Figure 2a,b and Figure S3 in Supporting Information show the SEM images (right-hand images) of the corresponding AFM ones. Films deposited at low oxygen pressures (PO 2 < 350 mTorr) show a rather flat morphology covered with small grains distributed homogeneously over the film surface. The latter corroborates the morphology observed by AFM and the increase of the roughness of the films. On the other hand, at high oxygen pressures (PO 2 ≥ 350 mTorr), the SEM images reveal a flat and smooth morphology, although some small grains randomly distributed over the film surface can be observed. As mentioned before, high pressure favors a smoother surface. The presence of these small grains at the film surface may be due to the presence of the NiO secondary phase, as previously detected by the XRD patterns (see Figure 1a,b). On the other hand, the existence of a nanostructured surface was also detected in the LNMO systems, and it was attributed to the formation of NiO x aggregates after high annealing temperatures (800-900 • C) [30].

Deposition Temperature
According to the results on the PNMO films grown at different oxygen pressures, good ferromagnetic properties were obtained for the samples deposited at intermediate oxygen pressures ranging from 350 to 400 mTorr O 2 (vide ultra). Therefore, in order to further optimize the deposition conditions, the PNMO thin films on the STO (001) substrates were grown at different substrate temperatures (ranging from 200 • C to 900 • C) by keeping a fixed oxygen pressure of 350 mTorr for 1 h. Figure S4a in Supporting Information shows a typical XRR curve of a PNMO film deposited at high temperature (T = 700 • C). The thickness dependence as a function of the growth temperature is shown Figure S4b in Supporting Information. It can be observed that the film thickness remains almost constant above 600 • C. The thickness values found via the XRR analysis for all the PNMO samples grown at different temperatures are very similar (~48 nm), with a growth rate of about 0.8 nm/min. In a particular case, the sample deposited at a substrate temperature of 200 • C was thicker (~54 nm). Figure 3a shows the XRD θ/2θ scans of the~48 nm-thick PNMO films grown under 350 mTorr O 2 at different temperatures. In addition, in all the XRD patterns, only the (00l) reflections of the PNMO phase were clearly observed in the samples grown at temperatures above 600 • C (see the dashed vertical red lines), indicating that high growth temperatures favor a good crystalline quality. For more details, Figure 3b shows the (002) reflection of both the STO substrate and the PNMO film. As the growth temperature increases, it can  Figure 3b). This shift indicates a progressive shrinking of the out-of-plane c lattice parameter (see Figure 3c) due to the tensile in-plane strain imposed by the substrate. The c lattice parameter of the films decreases from 3.866 Å to 3.851 Å as the growth temperature increases from 600 • C to 900 • C, as shown in Figure 3c.  Figure 3b). This shift indicates a progressive shrinking of the out-of-plane c lattice parameter (see Figure 3c) due to the tensile in-plane strain imposed by the substrate. The c lattice parameter of the films decreases from 3.866 Å to 3.851 Å as the growth temperature increases from 600 °C to 900 °C, as shown in Figure 3c.  Figure S1).
In addition to the PNMO reflections, the XRD patterns show the existence of a secondary NiO phase (see the asterisk and dashed blue line in Figure 3a,b), as previously detected in the samples deposited at different pressures ( Figure 1). Nevertheless, for the films grown/annealed in the 600-700 °C range, the (002) NiO reflection has a very low intensity. As the growth temperature increases, it can be observed that the NiO peak is slightly more intense and presents a slight shift towards lower angles (2θ ≈ 43.8° to 43.6° for 800-900 °C, respectively). As discussed before, this phase segregation may be of the same origin as seen in the LNMO thin films [30]. Therefore, according to the XRD patterns, the PNMO films deposited under 350 mTorr O2 at different temperatures grow epitaxially on the STO substrates, and with a good crystalline quality.
In order to determine the values of the in-plane (a) and out-of-plane (c) lattice parameters, a reciprocal space map (RSM) around the (−103) reflection was performed on an illustrative PNMO film with optimized growth conditions (grown/annealed at 800 °C under 350 mTorr O2). The RSM (Figure 4) shows that the peak position qx of the (−103) PNMO film reflection is slightly shifted towards a larger (absolute) value of qx from that of the STO substrate. The latter is indicating that the in-plane (a) cell parameter of the PNMO film is slightly smaller than the STO one (aSTO = 3.905 Å). This contrasts with the PNMO films grown on top of (LaAlO3)0.3(Sr2AlTaO6)0.7 (LSAT) [27], although this can be attributed to the smaller lattice mismatch introduced by LSAT (aLSAT = 3.869Å). Therefore, from the position of the (−103) PNMO film reflection, the in-plane lattice parameter (in pseudocubic notation) is found to be a = 3.877Å, which is closer to the lattice parameter of pseudocubic bulk PNMO (abulk = 3.871 Å), indicating a (partial) in-plane relaxation of the PNMO film. In addition, the out-of-plane lattice parameter of the PNMO film (c= 3.848 Å), determined from the qz value of the (−103) PNMO reflection, is smaller than that of the STO, in agreement with the θ/2θ scans shown in Figure 3.  Figure S1).
In addition to the PNMO reflections, the XRD patterns show the existence of a secondary NiO phase (see the asterisk and dashed blue line in Figure 3a,b), as previously detected in the samples deposited at different pressures ( Figure 1). Nevertheless, for the films grown/annealed in the 600-700 • C range, the (002) NiO reflection has a very low intensity. As the growth temperature increases, it can be observed that the NiO peak is slightly more intense and presents a slight shift towards lower angles (2θ ≈ 43.8 • to 43.6 • for 800-900 • C, respectively). As discussed before, this phase segregation may be of the same origin as seen in the LNMO thin films [30]. Therefore, according to the XRD patterns, the PNMO films deposited under 350 mTorr O 2 at different temperatures grow epitaxially on the STO substrates, and with a good crystalline quality.
In order to determine the values of the in-plane (a) and out-of-plane (c) lattice parameters, a reciprocal space map (RSM) around the (−103) reflection was performed on an illustrative PNMO film with optimized growth conditions (grown/annealed at 800 • C under 350 mTorr O 2 ). The RSM (Figure 4) shows that the peak position q x of the (−103) PNMO film reflection is slightly shifted towards a larger (absolute) value of q x from that of the STO substrate. The latter is indicating that the in-plane (a) cell parameter of the PNMO film is slightly smaller than the STO one (a STO = 3.905 Å). This contrasts with the PNMO films grown on top of (LaAlO 3 ) 0.3 (Sr 2 AlTaO 6 ) 0.7 (LSAT) [27], although this can be attributed to the smaller lattice mismatch introduced by LSAT (a LSAT = 3.869Å). Therefore, from the position of the (−103) PNMO film reflection, the in-plane lattice parameter (in pseudocubic notation) is found to be a = 3.877Å, which is closer to the lattice parameter of pseudocubic bulk PNMO (a bulk = 3.871 Å), indicating a (partial) in-plane relaxation of the PNMO film. In addition, the out-of-plane lattice parameter of the PNMO film (c= 3.848 Å), determined from the q z value of the (−103) PNMO reflection, is smaller than that of the STO, in agreement with the θ/2θ scans shown in Figure 3.  The AFM and SEM images of the PNMO films for the different growth temperatures are shown in Figure 5a,b and Figure S5 in Supporting Information. It can be observed that for the PNMO films grown at low temperatures, from 200 • C to 700 • C, the AFM and SEM images (left-hand images) show a flat surface without the presence of grains. Moreover, the roughness values of the films obtained from the AFM data are below 0.4 nm (see Figure 5c), indicating that the surface roughness is similar to that of the steps-terraces morphology of the underlying substrates. At higher growth temperatures (T ≥ 800 • C), the films show a granular morphology over the entire surface (see Figure S5 in Supporting Information). In particular, the sample grown at 900 • C has a surface morphology with a higher grain density (according to the SEM image in Figure S5 in Supporting Information), which is also reflected in an increase of the roughness (rms ≈ 0.7 nm see Figure 5c). Nevertheless, a low roughness value (rms < 1 nm) can be appreciated in all the PNMO films grown at 350 mTorr O 2 . In addition, the XRD patterns showed the existence of a secondary NiO phase (see Figure 3a,b). The NiO peak becomes more intense at high growth/annealing temperatures (800-900 • C), which could be related to the formation of NiO x segregates at the film surface at high annealing temperatures (800-900 • C), as observed previously in the LNMO films [30]. For the films grown/annealed in the 200-700 • C range, the secondary NiO phase was not clearly detected by XRD, as corroborated by the AFM and SEM images.
Supporting Information), which is also reflected in an increase of the roughness (rms ≈ 0.7 nm see Figure 5c). Nevertheless, a low roughness value (rms < 1 nm) can be appreciated in all the PNMO films grown at 350 mTorr O2. In addition, the XRD patterns showed the existence of a secondary NiO phase (see Figure 3a,b). The NiO peak becomes more intense at high growth/annealing temperatures (800-900 °C), which could be related to the formation of NiOx segregates at the film surface at high annealing temperatures (800-900 °C), as observed previously in the LNMO films [30]. For the films grown/annealed in the 200-700 °C range, the secondary NiO phase was not clearly detected by XRD, as corroborated by the AFM and SEM images.

Chemical Composition
The chemical composition and the stoichiometry ratio of the PNMO thin films grown on the STO substrates under different oxygen pressures and temperatures have been analyzed by EPMA. Measurements were done at 10 different points of each sample in order to estimate the error bars of the measurement. Figure 6a,b show the stoichiometry ratio obtained.

Chemical Composition
The chemical composition and the stoichiometry ratio of the PNMO thin films grown on the STO substrates under different oxygen pressures and temperatures have been analyzed by EPMA. Measurements were done at 10 different points of each sample in order to estimate the error bars of the measurement. Figure 6a,b show the stoichiometry ratio obtained. In Figure 6a, it can be appreciated that the stoichiometry ratio of the PNMO films (grown at 850 °C) shows a strong dependence on the oxygen pressure PO2. Samples grown at low oxygen pressures (ranging from 100 to 300 mTorr) present a Pr/(Ni+Mn) atomic slightly above 1 (see black symbols), while the Ni/Mn ratio is below 1 (see red symbols). One can conclude that these samples are Ni-deficient. Contrarily, in samples grown at high oxygen pressures (PO2 ≥ 350 mTorr), the Pr:(Ni+Mn) ratio is clearly below 1, while the Ni/Mn ratio is close to 1, indicating that the films have a certain degree of Pr deficiency. On the other hand, the error in determining the O content is larger than for the other elements due to the presence of oxygen in both the film and substrate. In this case, the oxygen concentration is not computed due to the presence of a large O background signal from the substrate, which is difficult to separate from the one coming from the film. Therefore, stoichiometry is not considered for the oxygen and the composition of the samples that are expressed as Pr2−δNi1−xMn1+xO6−y when Pr/(Ni+Mn) ≤ 1, and Pr2Ni1−xMn1−zO6−y when Pr/(Ni+Mn) ≥ 1, as the presence of oxygen vacancies (y>0) cannot be ruled out. In this regard, the samples grown at 100 mTorr and 200 mTorr O2 show a Ni/Mn ratio of about 0.58 (Pr2.0Ni0.7Mn1.2O6−y) and 0.62 (Pr2.0Ni0.7Mn1.1O6−y), respectively, i.e., with Ni deficiency which thus affects the ferromagnetic properties (as discussed in the next subsection). At high oxygen pressures (PO2 ≥ 350 mTorr), the Pr:(Ni+Mn) ratio is below 1: films grown at 350 mTorr and 400 mTorr O2 show a Pr:(Ni+Mn) ratio of ~0.83 In Figure 6a, it can be appreciated that the stoichiometry ratio of the PNMO films (grown at 850 • C) shows a strong dependence on the oxygen pressure PO 2 . Samples grown at low oxygen pressures (ranging from 100 to 300 mTorr) present a Pr/(Ni+Mn) atomic slightly above 1 (see black symbols), while the Ni/Mn ratio is below 1 (see red symbols). One can conclude that these samples are Ni-deficient. Contrarily, in samples grown at high oxygen pressures (PO 2 ≥ 350 mTorr), the Pr:(Ni+Mn) ratio is clearly below 1, while the Ni/Mn ratio is close to 1, indicating that the films have a certain degree of Pr deficiency. On the other hand, the error in determining the O content is larger than for the other elements due to the presence of oxygen in both the film and substrate. In this case, the oxygen concentration is not computed due to the presence of a large O background signal from the substrate, which is difficult to separate from the one coming from the film. Therefore, stoichiometry is not considered for the oxygen and the composition of the samples that are expressed as Pr 2−δ Ni 1−x Mn 1+x O 6−y when Pr/(Ni+Mn) ≤ 1, and Pr 2 Ni 1−x Mn 1−z O 6−y when Pr/(Ni+Mn) ≥ 1, as the presence of oxygen vacancies (y > 0) cannot be ruled out. In this regard, the samples grown at 100 mTorr and 200 mTorr O 2 show a Ni/Mn ratio of about 0.58 (Pr 2.0 Ni 0.7 Mn 1.2 O 6−y ) and 0.62 (Pr 2.0 Ni 0.7 Mn 1.1 O 6−y ), respectively, i.e., with Ni deficiency which thus affects the ferromagnetic properties (as discussed in the next subsection). At high oxygen pressures (PO 2 ≥ 350 mTorr), the Pr:(Ni+Mn) ratio is below 1: films grown at 350 mTorr and 400 mTorr O 2 show a Pr:(Ni+Mn) ratio of~0.83 (Pr 1.7 Ni 0.9 Mn 1.1 O 6−y ) and~0.73 (Pr 1.5 Ni 1.0 Mn 1.0 O 6−y ), respectively. The slight deficiency of Pr in the films when the oxygen pressure increases could be attributed to the fact of a greater scattering of plasma species, being that this scattering more notable for the lighter elements. Therefore, a deficient flow of adatoms on the film surface could render the formation of nonstoichiometric Pr/(Ni+Mn) < 1.
Finally, Figure 6b shows the atomic ratio of the PNMO films deposited at different growth temperatures for PO 2 = 350 mTorr. It is found that all samples present a Pr/(Ni + Mn) atomic ratio clearly below 1, while Ni/Mn is near 1, indicating that the films show a Pr deficiency. At the same time, it can be also observed that by increasing the growth temperature (T ≥ 700 • C), the amount of Pr slightly increases but is still below the expected stoichiometry. For instance, the film grown/annealed at 800 • C shows Pr:(Ni + Mn) ≈ 0.86 and Ni/Mn ≈ 0.86 (Pr 1.7 Ni 0.9 Mn 1.1 O 6−y ).

Magnetic and Transport Properties
The magnetization as a function of temperature was measured by applying a 5 kOe magnetic field in the plane of the samples after a field cooling process (at the same field and orientation). The resulting curves are shown in Figure 7a,d. Curie temperatures were estimated from the inflection point of the M(T) curves and are plotted in Figure 7b,e. As can be seen, a quite significant magnetic moment appears in all films grown at 850 • C (Figure 7a), although the T C (Figure 7b) presents a relevant dependence on the oxygen pressure between 100 mTorr and 350 mTorr. In this region, a fast increase of the Curie point from 118 K to about 200 K takes place. At 350 mTorr (and up to 750 mTorr), T C reaches a plateau and remains almost constant around 200 K. In addition, in this growth pressure range, the M(T) curves do not show signs of any significant secondary magnetic transition at low temperature (in contrast with previously reported films grown by PLD [26,27]). These results clearly demonstrate that the ordering of Ni 2+ and Mn 4+ ions is strongly affected by the oxygen pressure during deposition. The value of the FM ordering temperature in the RNMO structure is determined by the magnitude of the superexchange spin-transfer integral [32,33]. The low-T C phase of RNMO is attributed to trivalent oxidation states, Ni 3+ and Mn 3+ , while the high-T C phase has been shown to be characteristic of the atomically ordered Ni 2+ and Mn 4+ [34]. It is worth mentioning that other double perovskites, such as La 2 MnCoO 6 , present an important dependence of the magnetic properties on oxygen content in both bulk and thin film forms [35,36]. Oxygen vacancy concentration increases when the films are deposited at low oxygen pressure, which breaks the Ni 2+ -O-Mn 4+ superexchange ferromagnetic interaction and leads to higher spin disorder, hence decreasing the T C .
M(H) hysteresis loops measured at 10 K with H applied in-plane corresponding to two samples grown under oxygen pressure of 350 and 400 mTorr O 2 with T C values about 200 K are shown in Figure 7c. The shape of the M(H) loop is practically identical for both samples. Consequently, a similar coercive field (H C ≈ 500 Oe) and the same saturation magnetization (M s ≈ 5 µ B /f.u) were extracted for them. Magnetization remanence M r slightly decreases by increasing the oxygen pressure from M r ≈ 2.7 µ B /f.u to 2.6 µ B /f.u, for 350 and 400 mTorr O 2 , respectively. The estimated values of the saturation magnetization equal to the optimal spin-only value of 5 µ B /f.u [6,19,22] is encountered in fully ordered Pr 2 NiMnO 6 systems. Thus, the PNMO films grown at 350 and 400 mTorr O 2 , in spite of the Pr deficiency shown by EPMA (indicating compositions Pr 1.7 Ni 0.9 Mn 1.1 O 6−y and Pr 1.5 Ni 1.0 Mn 1.0 O 6−y , respectively), present very optimal magnetic properties with quite high T C and low temperature saturation magnetization nearby to the maximum expected for Ni 2+ and Mn 4+ . These magnetic properties are similar to those previously reported by M. P. Singh et al., for PNMO films prepared by PLD [26]. They report good FM properties (i.e., T C ≈ 214 K and M s ≈ 4.35 µ B /f.u at 10 K) using higher oxygen pressures, i.e., 900 mTorr O 2 and growth temperatures of 800 • C. Nevertheless, the presence of an additional magnetic transition T C2 ≈ 126 K was assigned to cation disordering and the coexistence of multiple oxidation states of Ni and Mn. The oxygen pressure used during the deposition of the PNMO thin films plays an important role to stabilize the distinct valence states of Ni and Mn cations in the films, which subsequently affects the FM properties. In our films grown/annealed at 850 • C under different oxygen pressures, the optimal T C and M S (very close to the optimal values [16,21,26]) have been obtained only at oxygen pressures in the 350-400 mTorr O 2 range. The analysis of the dependence on the substrate temperature, as mentioned, has been conducted at the optimal oxygen pressure (350 mTorr O 2 ). The magnetic properties of the PNMO thin films have been evaluated on films prepared at different growth temperatures (200 • C ≤ T ≤ 900 • C) under a fixed oxygen pressure of 350 mTorr O 2 . In Figure 7d, all M(T) curves are characterized by a single magnetic transition. The evolution of T C as a function of the growth temperature is shown in Figure 7e. Samples grown/annealed at temperatures below or equal to 600 • C show very deficient magnetic properties, with T C ≤ 86 K, and when the growth temperature increases to 800 • C, T C rises up to~210 K. At higher temperatures, from 850 • C on, the ferromagnetic transition temperature slightly decreases again. According to these results, the reduced T C in samples grown/annealed below 700 • C could be attributed to the fact that the growth temperature may not be sufficient to promote ordering of Ni and Mn cations. Ordering processes take place by a thermally activated diffusion of cations. In this regard, higher temperatures favor growth kinetics and tend to increase ordering [37,38]. Therefore, growing/annealing at high temperatures (T ≥ 700 • C) promotes Ni/Mn cation ordering and consequently an increase of T C to values very close to that of the bulk one (i.e., T C ≈228 K and Ms ≈ 4.95 µ B /f.u at 50 kOe field) [21]. At the same time, the absence of a secondary transition at low temperatures in our films indicates a low concentration of antisite disorder.
M(H) hysteresis loops corresponding to two PNMO samples with good Curie temperature, grown/annealed at 800 • C (T C~2 10 K) and 850 • C (T C~2 01 K), measured at 10 K with a magnetic field applied in-plane, are shown in Figure 7f. For both samples, the hysteretic M(H) behavior can be clearly observed, confirming the ferromagnetic character of the PNMO thin films. The coercive field H C was found to be H C = 565 Oe and 482 Oe for samples deposited at 800 • C and 850 • C, respectively. At the same time, the remnant magnetization M r increases from M r ≈ 2.3 µ B /f.u to 2.7 µ B /f.u and the saturation magnetization increases from M s ≈ 4.5 µ B /f.u. to 5 µ B /f.u, while increasing the growth temperature from 800 • C to 850 • C. On the other hand, in our results, although the T C is slightly higher in the sample grown at 800 • C, this film showed a lower saturation magnetization value than that of the film grown at 850 • C. This reduction of the magnetization could be caused by the formation of antiphase boundaries and alternatively some antisite defects (i.e., Ni-O-Ni or Mn-O-Mn) [24,39]. Therefore, films with a T C and Ms very close to the optimal values [16,21,26] have been obtained only at growth temperatures in the 800-850 • C range under an oxygen pressure of 350 mTorr, in spite of their slight Pr deficiency.
The temperature dependence of the resistivity ρ(T) measured in the range 150 K ≤ T ≤ 300 K (at zero magnetic field) of a PNMO film with optimized growth conditions (i.e., grown/annealed at 800 • C under 350 mTorr O 2 ), displaying good ferromagnetic properties (T C ≈ 210 K and Ms ≈ 4.5 µ B /f.u. at 10 K), is shown in Figure 8a. The measurement shows that the resistivity increases with decreasing the temperature, thus exhibiting an insulating behavior in the whole measured temperature range. At lower temperatures (T < 150 K), ρ(T) was not measurable due to the high sample resistance. The dependence of ρ(T) can be well fitted by using the thermal activation mechanism, described by the Arrhenius law (see Figure 8b) ρ(T) = ρ 0 exp(E 0 /K B T), where ρ 0 is a constant, E 0 is the required activation energy for the conduction process generally at high temperatures, and K B is the Boltzmann constant [40,41]. This model can explain the thermally generated charge-transport behavior and activation type mechanism in the insulating state [42]. The fitting of the Arrhenius model was found in a temperature range from 300 K to 195 K. The extracted value of the activation energy E 0 for our PNMO film was found to be E 0 ≈ 0.18 ± 0.1 eV, which is comparable to that of the cation-ordered RNMO compounds [40,42,43].

Conclusions
In conclusion, we have achieved the epitaxial growth of double perovskite PNMO thin films on (001) STO substrates by the RF magnetron sputtering technique. It is shown that ferromagnetic ordering is significantly influenced by both the oxygen pressure and the growth/annealing temperature. The optimization of the growth conditions has revealed that only under a narrow window of deposition conditions (an oxygen pressure of 350-400 mTorr and growth/annealing temperatures around 800-850 • C), PNMO thin films with good ferromagnetic properties (i.e., T C ≥ 210 K and Ms ≥ 4.5 µ B /f.u. at 10K very close to the bulk value) and insulating behavior can be prepared by using AC magnetron sputtering. High annealing temperatures (800-900 • C) promote the appearance of a nanostructured surface, attributed to the formation of NiO x segregates, as previously observed in LNMO systems. Additionally, the XRD patterns showed the existence of a secondary NiO phase. However, although secondary phases are formed (NiO) together with the PNMO phase, films grow epitaxially on the (001) STO substrates. Regarding the stoichiometry of the films (measured by EPMA), it is found that the optimization of the growth parameters does not imply obtaining stoichiometric samples of the expected Pr 2 NiMnO 6 compound. Particularly, the PNMO films grown/annealed at different growth temperatures under a high oxygen pressure (PO 2 ≥ 350 mTorr) show a Ni/Mn ratio close to 1 and a Pr:(Ni + Mn) ratio with a certain degree of Pr deficiency, and in spite of this Pr deficiency, the samples exhibit good ferromagnetic and electrical transport properties.